PROCESSING AND PROPERTIES OF A DUAL PHASE PM
STEEL
Chris Schade & Tom Murphy
Hoeganaes Corporation
Cinnaminson, NJ 08077
Alan Lawley & Roger Doherty
Drexel University
Philadelphia, PA 19104
ABSTRACT
The development of materials utilizing lower alloy contents is essential to the cost
competiveness of powder metallurgy (PM). Alloys systems that contain low levels of
alloying elements such as nickel and molybdenum, yet provide adequate strength and
ductility, can be developed utilizing strengthening mechanisms such as precipitation
hardening and microalloying. This study details the processing, microstructures and
properties of a low alloy PM steel in which these alloying techniques are used with the
ultimate goal of producing a dual-phase microstructure coupled with precipitation
hardening. Processing of these alloys is detailed and attendant microstructures and
mechanical properties compared with those of conventional PM steels.
INTRODUCTION
There is a continuing need to develop PM steels with higher strength, while maintaining
adequate levels of ductility and impact toughness. Although advances have been made in
high density processing, such as warm compaction, improved lubrication, and high
density powders, there remains a need to develop alloy systems which, in themselves,
improve the competitiveness of PM in relation to wrought alloys. In these systems it is
desirable to use as little alloy content as possible. High-strength low alloy (HSLA) steels
or microalloyed steels, generally give superior mechanical properties utilizing low alloy
levels coupled with thermo-mechanical processing. Dual-phase steels are considered a
subclass of HSLA steels and exhibit a microstructure consisting of a hard phase
(primarily martensite and/or bainite) in a matrix of ferrite. Due to their composite
microstructures, dual phase steels exhibit excellent mechanical properties with tensile
strength generally dependent primarily on the volume fraction of martensite. The authors
have recently developed a dual phase precipitation hardened stainless steel in which this
combination of strengthening mechanisms has led to a material with improved strength
and ductility1.
To explore whether or not this technique can be applied to conventional low alloy PM
steels a development study was undertaken in which the precipitation of copper was
studied in iron-copper systems typically used in low alloy PM steels. Because the
precipitation temperature for copper is relatively high (450 to 600 oC (850 to 1100 oF)),
the addition of alloying elements which tend to stabilize the carbides was examined.
Since copper stabilizes austenite (which transforms to martensite on cooling), elements
that stabilize ferrite such as chromium, molybdenum and silicon were studied with the
expectation that a dual phase microstructure could be developed. In addition,
microalloying elements such as niobium, titanium and vanadium were evaluated in an
attempt to co-precipitate carbides and nitrides along with the copper. This may allow for
the use of lower levels of alloying elements, including carbon. This approach is attractive
to the PM industry since low carbon contents yield higher green and sintered densities. It
is also possible that the loss of strength due to the lower carbon level can be offset by
precipitation.
Precipitation Hardening
Strengthening, as a result of precipitation hardening, takes place in three steps:2
(1) Solution treatment, in which the alloy is heated to a relatively high temperature,
allows any precipitates or alloying elements to form a supersaturated solid
solution. Typical solution treatment temperatures for steels are in the range of
982 oC to 1066 oC (1800 oF to 1950 oF).
(2) Quenching, in which the solution treated alloy is cooled to create a supersaturated
solid solution. The cooling can be achieved using air, water or oil. In general, the
faster the cooling rate the finer the grain size which can lead to improved
mechanical properties. Regardless of the method of cooling, the cooling rate must
be sufficiently rapid to create a supersaturated solid solution.
(3) Precipitation or age hardening, in which the quenched alloy is heated to an
intermediate temperature or held at room temperature for a period of time.
During aging, the supersaturated solid solution decomposes and the alloying
elements form small precipitate clusters. The precipitates hinder the movement of
dislocations and consequently the metal resists deformation and becomes harder
and stronger.
The age hardening response of copper-bearing high strength low alloy and multiphase
steels has been shown to occur at temperatures ranging from 450 to 600 oC (850 to 1100
o
F) and has been shown to increase strength and hardness by 15 to 20 percent 3-5.
Since there is widespread use of the iron-copper system in PM components, a study of
aging in the iron-copper system to enhance strength is a logical step.
However there are a number of challenges to this objective, as illustrated by Figure 1
which shows the change in transverse rupture strength with aging temperature for
o
Aging Temperature ( C)
93
204
315
427
538
170
1170
0.8 w/o
160
1101
0.6 w/o
140
963
0.4 w/o
130
894
120
TRS (MPa)
1032
3
TRS (x10 psi)
150
826
0.2 w/o
110
757
0.0 w/o
100
688
90
0
200
400
600
800
1000
619
1200
o
Aging Temperature ( F)
Figure 1: TR strength of Fe-3 w/o Cu at various carbon levels versus aging temperature
several carbon levels. A classical aging response can be seen in the carbon free alloy,
namely an increase in strength as the aging temperature is increased, followed by a
decrease at higher temperatures as the copper precipitates coarsen. As the carbon level is
increased there is a gain in the as sintered strength (“0” aging time) but the aging
response is less pronounced. During the aging process two important microstructural
features prevail: the copper precipitates and increases the strength and hardness of the
PM steel, but concurrently the high temperatures necessary for this reaction results in
coarsening of the carbides, offsetting any benefit from precipitation hardening. Therefore
there are two competing processes; precipitation of copper and tempering of the carbides.
Tempering
Since the effect of carbon on hardness and strength decreases markedly with increasing
tempering temperature, a means to stabilize the carbides at higher temperatures is needed
if the carbides are to contribute to the strength of an alloy aged at 538 oC (1000 oF).
Although there is published information on the effect of specific alloying elements on the
hardness of tempered martensite, it is based on medium carbon steels and does not
account fully for the effects of secondary hardening and precipitation.6-7 The relationship
between carbon content and apparent hardness in as quenched martensite is shown as a
function of tempering temperature in Figure 2. Maximum hardness is acheived in the
quenched martensite and decreases with increasing tempering temperature. For Fe-0.4
w/o C the hardness drops from approximatley 700 HV (as quenched) to 210 HV when
tempered at 538 oC (1000 oF). This is the temperature regime at which copper
precipitation is maximized. The primary purpose of tempering is to impart toughness to
the alloy. This occurs due to the relief of internal stresses caused by the formation of
martensite and the precipitation and spherodization of carbides. Alloying elements have
a direct effect on the latter. Elements such as manganese, nickel, and silicon have only a
small effect since they do not form carbides.
Figure 2. Hardness of tempered martensite in iron-carbon alloys6.
Elements such as chromium, molybdenum and vanadium are known to enhance hardness
at higher tempering temperatures by forming alloy carbides. The effectiveness of these
elements depends on their solubility in austenite and the nature of the carbide formed.
These alloying elements also retard the coalescence of carbides, leading to temper
resistance. Figure 3 shows the effect of various alloying elements (and concentrations)
on the hardness of the martensite in a Fe-0.2 w/o C alloy. Tempering was performed at
538 oC (1000 oF) for 1 h. At low concentration levels of alloying molybdenum and
vanadium result in effective temper resistance, particularly at 538 oC (1000 oF) since this
is the temperature at which the precipitation of copper takes place.
Figure 3. Effect of alloying elements on the hardness of martensite 6.
The objective of the present study was to develop and evaluate an iron-copper–carbon
PM alloy that could be aged at 538 oC (1000 oF) while maintaining significant temper
resistance in order to retain strength and apparent hardness. This outcome provides the
necessary framework to produce an alloy at carbon levels in the range of 0.4 to 0.6 w/o
with a multiphase microstructure.
ALLOY PREPARATION AND TESTING
The powders used in this study were produced in both the prealloy and admixed
conditions. Where pertinent, the processing steps and powder types are detailed in the
‘discussion’ section
The powders were mixed with Acrawax C lubricant and graphite. Samples for transverse
rupture (TR) and tensile testing were compacted uniaxially at pressures of 690 MPa and
730 MPa (50 tsi and 53 tsi). The test pieces were sintered in a high temperature Abbott
continuous-belt furnace at a temperature of 1120 °C (2050 °F) for 30 min in a mixed
atmosphere of 90 v/o nitrogen and 10 v/o hydrogen.
Prior to mechanical testing, green and sintered density, dimensional change (DC), and
apparent hardness, were determined on the tensile and TR samples. Five tensile
specimens and five TR specimens were evaluated for each composition. The densities of
the green and sintered steels were determined in accordance with MPIF Standard 42.
Tensile testing followed MPIF Standard 10 and apparent hardness measurements were
conducted on tensile and TR specimens, in accordance with MPIF Standard 43.
Specimens of the test materials were prepared using standard metallographic procedures
and were examined by optical microscopy in the polished and etched (1 v/o nital / 4 w/o
picral.) conditions.
RESULTS AND DISCUSSION
Temper Resistance as a Function of Alloying Element
In order to establish the effects of chromium, molybdenum and silicon on tempering, a
Fe-3 w/o Cu base alloy was used. The copper added was Acupowder 8081 and admixed
with 0.60 w/o Acrawax C with additions of graphite (from 0.20 to0.80 w/o). Where
noted in the ‘discussion’, elements were added either as ferroalloy additives or prealloy
base powders were used to evaluate temper resistance. TR specimens were pressed and
sintered to a density of 7.0 g/cm3. The TR strength and apparent hardness were used as
criteria to assess the effectiveness of the alloy elements in relation to temper resistance.
Conventional Alloying Elements
Chromium
The effect of chromium was studied by adding high carbon ferrochromium to achieve
chromium levels of 0.25, 0.50 and 1.00 w/o (at 0.40 w/o carbon). The tempering response
of these alloys is compared with that of a Fe-3.0 w/o Cu base alloy in Figure 4. The
effect of chromium on temper resistance is minimal. Although chromium is a strong
carbide former, the carbide forms at temperatures ranging from 204 to 427 oC (400 to 800
o
F). At 538 oC (1000 oF) the precipitation of copper occurs. At this temperature the
carbides may already be coalescing. Based on Figure 4, significant temper resistance
starts to occur at 0.50 w/o Cr.
o
o
Tempering Temperature ( C)
Tempering Temperature ( C)
204
315
427
538
93
3
TRS (x10 psi)
1200
170
0.25 w/o Chromium
1120
160
1040
150
0 w/o Chromium
140
960
130
TRS (MPa)
1280
180
315
427
538
0.75 w/o Chromium
190
0.50 w/o Chromium
204
60
1360
0.75 w/o Chromium
Apparent Hardness (HRA)
93
200
58
56
0.50 w/o Chromium
54
0.25 w/o Chromium
52
50
48
0 w/o Chromium
46
44
0
200
400
600
800
1000
o
Tempering Temperature ( F)
(a)
1200
0
200
400
600
800
1000
o
Tempering Temperature ( F)
(b)
Figure 4. Temper resistance of Fe-3 w/o Cu with various levels of chromium: (a) TR strength and (b)
apparent hardness. Nominal density 7.0 g/cm3.
1200
Molybdenum
The effect of molybdenum was studied by adding 3.0 w/o Acupowder 8081 copper and
0.40 w/o graphite to Ancorstreel 50HP, Ancorsteel 85HP and Ancorsteel 150HP which
have nominal molydenum levels of 0.50, 0.85 and 1.50 w/o respectively. The tempering
response of these materials is compared with that of a molybdenum-free Fe-3.0 w/o Cuiron base alloy in Figure 5.
o
o
Tempering Temperature ( C)
Tempering Temperature ( C)
93
204
315
427
538
93
204
315
427
538
58
180
1.50 w/o Molybdenum
1120
160
1080
0.85 w/o Molybdenum
1040
150
1000
0 w/o Molybdenum
140
960
TRS (MPa)
3
TRS (x10 psi)
1160
Apparent Hardness (HRA)
1200
0.50 w/o Molybdenum
170
56
1.50 w/o Molybdenum
54
52
0.85 w/o Molybdenum
50
0.50 w/o Molybdenum
48
0 w/o Molybdenum
46
920
130
44
0
200
400
600
800
1000
1200
o
Tempering Temperature ( F)
(a)
0
200
400
600
800
1000
o
Tempering Temperature ( F)
(b)
Figure 5. Temper resistance of Fe-3 w/o Cu with various levels of molybdenum: (a) TR strength and (b)
apparent hardness. Nominal density 7.0 g/cm3.
Molybdenum is a strong carbide former and results in higher strength and apparent
hardness, compared with molybdenum-free Fe-3.0 w/o Cu. At temperatures below 316
o
C (600 oF) molybdenum had no effect but at temperatures > 427 oC (800 oF) increases in
strength and apparent hardness were seen. In wrought grades, molybdenum is known to
be an effective alloying addition at tempering temperatures > 538 oC (1000 oF) because it
partitions to the carbide phase at elevated temperatures and prevents coarsening of the
carbides. The small observed differences in strength may be due to the small specimen
size; larger differences in strength would be expected on larger parts.
Silicon
Silicon is a solid solution strengthener and is also effective in refining carbides during
tempering. There is evidence that silicon delays the ε-carbide to cementite
transformation which normally occurs at 316 oC (600 oF). It has been shown that,
although silicon is not soluble in the carbide, if present in significant quantities it hinders
the diffusion of carbon and slows the coarsening rate of the carbides. Thus higher
tempering temperatures are required to soften the alloy. Newer steels that have been
developed for springs take advantage of the solid solution strengthening of silicon,
1200
coupled with its temper resistance in the production of higher strength – sag resistant
components.9
o
o
Tempering Temperature ( C)
Tempering Temperature ( C)
93
204
315
427
93
538
204
427
315
538
52.5
174
1190
1160
3
TRS (x10 psi)
1170
170
168
1150
0.40 w/o Silicon
1140
166
1130
164
TRS (MPa)
1180
Apparent Hardness (HRA)
0.20 w/o Silicon
172
52
51.5
0.40 w/o Silicon
51
50.5
0.20 w/o Silicon
50
1120
0.60 w/o Silicon
0.60 w/o Silicon
49.5
162
0
200
400
600
800
1000
1200
o
Tempering Temperature ( F)
(a)
0
200
400
600
800
1000
1200
Tempering Temperature (F)
(b)
Figure 6. Temper resistance of Fe-3 w/o Cu with various levels of silicon: (a) TR strength and (b)
apparent hardness. Nominal density 7.0 g/cm3.
The results for an iron-3w/o copper with various amount of silicon are shown in Figure 6.
At lower levels of silicon (0.20 w/o) there is no significant change in transverse rupture
strength and the hardness decreases slightly with tempering temperatures. As the silicon
level increases (0.40 and 0.60 w/o) the transverse rupture strength increases slightly at
aging temperatures above 600 oF (316 oC), while the hardness does not significantly
change. At all three silicon levels there is a slight drop in transverse rupture strength
when tempered below this temperature.
Microalloying Elements
Since microalloying elements are normally added in amounts < 0.2 w/o, and and it is
difficult to achieve solid solution alloying during conventional sintering, additions were
made by prealloying. Mechanical property results indicated that 3 w/o copper in the
alloy was not necessary to maximize properties; thus 2 w/o copper was prealloyed into
the alloy. The following section details the effect of selected microalloying prealloy
additions to a 2 w/o copper prealloy.
Vanadium
Because vanadium carbide forms in steels at relatively small amounts (Figure 3), only
one level of vanadium was examined. Figure 7 shows the TR strength and apparent
hardness as a function of tempering temperature. Although the increase in apparent
hardness was small, the TR strength increased with increasing tempering temperatures at
427 oC (800 oF) and only decreased slightly above this temperature. The effect of
vanadium reflects the formation of an alloy carbide (probably V4C3 or VC) that replaces
cementite at high temperatures and persists as a fine precipitate that is well dispersed in
the matrix. The necessary requirement in the use of vanadium is that it goes into solid
solution at temperatures ranging from 1150 to 1250 oC (2100 to 2300 oF). By using a
ferroalloy addition, the full potency of the vanadium may not have been realized since
complete dissolution of the ferrovanadium particles may not have taken place during
sintering. Vanadium also has a high affinity for nitrogen which can lead to the
precipitation of vanadium carbonitride [V(C,N)]. This may be an advantage for PM
applications where the sintering atmosphere contains nitrogen.
o
o
Tempering Temperature ( C)
93
204
315
427
538
0.60 w/o Carbon
125
840
120
800
0.40 w/o Carbon
115
760
110
Apparent Hardness (HRA)
880
TRS (MPa)
920
130
204
315
427
538
55
0.80 w/o Carbon
135
3
93
960
140
TRS (x10 psi)
Tempering Temperature ( C)
0.80 w/o Carbon
50
0.60 w/o Carbon
45
0.40 w/o Carbon
40
0.20 w/o Carbon
0.20 w/o Carbon
105
35
0
200
400
600
800
1000
1200
o
Tempering Temperature ( F)
(a)
0
200
400
600
800
1000
1200
o
Tempering Temperature ( F)
(b)
Figure 7. Temper resistance of Fe-2 w/o Cu with vanadium at various levels of carbon: (a) TR strength
and (b) apparent hardness. Nominal density 7.0 g/cm3.
Niobium
Niobium results in significant strengthening via the precipitation of a hard and stable
carbide, nitride or carbonitride, NbC(N).10 The carbide-forming tendency of niobium is
stronger than that of vanadium and studies on wrought steel grades have shown that about
one-half the amount of niobium is needed compared with vanadium to produce the same
strength. In addition, the carbonitrides of niobium are more stable at higher temperatures
than those of vanadium. Niobium has a negative effect on the hardness of martensite
because of its ability to effectively remove carbon from solution and is generally not used
in low alloy high strength steels. Figure 8 shows the effect of niobium on TR strength and
apparent hardness for various tempering temperatures.
o
o
Tempering Temperature ( C)
204
93
427
315
Tempering Temperature ( C)
538
93
140
960
880
3
840
0.80 w/o Carbon
120
800
0.60 w/o Carbon
110
760
TRS (MPa)
TRS (x10 psi)
Apparent Hardness (HRA)
920
130
720
0.40 w/o Carbon
100
680
0.20 w/o Carbon
204
315
427
538
55
50
0.80 w/o Carbon
45
0.60 w/o Carbon
40
0.40 w/o Carbon
0.20 w/o Carbon
640
90
35
0
200
400
600
800
1000
1200
0
o
200
400
600
800
1000
1200
o
Tempering Temperature ( F)
Tempering Temperature ( F)
(a)
(b)
Figure 8. Temper resistance of Fe-2 w/o alloy with niobium at various levels of carbon: (a) TR strength
and (b) apparent hardness. Nominal density 7.0 g/cm3.
Niobium/Vanadium
The use of more than one carbide forming element makes it difficult to predict details of
the precipitation process. In general, if equilibrium is attained, the carbide phase that is
thermodynamically stable will predominate. However, commonly used tempering
temperatures in the range 500 to 600 oC of (932 to 1150 oF) do not allow for
thermodynamic equilibrium. The precipitates that form can contain more than one of the
elements or they can form a separate fine dispersion which is resistant to overaging. The
result of a combination of niobium and vanadium is shown in Figure 9.
o
o
Tempering Temperature ( C)
93
204
315
427
Tempering Temperature ( C)
538
93
204
315
427
538
55
160
1040
0.80 w/o Carbon
1000
960
140
920
0.60 w/o Carbon
130
880
0.40 w/o Carbon
840
120
800
TRS (MPa)
3
TRS (x10 psi)
150
Apparent Hardness (HRA)
1080
0.80 w/o Carbon
50
45
0.60 w/o Carbon
0.40 w/o Carbon
40
0.20 w/o Carbon
0.20 w/o Carbon
760
110
0
200
400
600
800
1000
Tempering Temperature (F)
(a)
1200
35
0
200
400
600
800
1000
1200
o
Tempering Temperature ( F)
(b)
Figure 9. Temper resistance of Fe-2 w/o Cu with niobium and vanadium with various levels of carbon: (a)
TR strength and (b) apparent hardness. Nominal density 7.0 g/cm3.
Mechancial Properties
Since the use of microalloys was shown to enhance temper resistance, the mechanical
properties of the vanadium steel was examined in the aged and heat treated conditions.
The use of vanadium allows for the strength and apparent hardness levels to be
maintained through the tempering cycle. Table I summarizes mechanical properties of
Fe-2 w/o Cu, with and without vanadium in the as sintered, aged and heat treated
conditions. In the aged condition the tensile properties of the two alloys are the same.
After aging 1 h at 538 oC (1000 oF) the tensile stength of Fe-2 w/o Cu with vanadium
shows an improvement in the yield strength, ultimate tensile strength, and elongation. In
the heat treated condition the strength of the alloy with vanadium exceeds that of
vanadium-free Fe-2 w/o Cu by 20 %.
Table I: Mechanical Properties of Fe-Cu and Fe-Cu-V PM alloys.
Alloy
Sintered
Density
Apparent
Hardness
(g/cm3)
7.00
6.98
(HRA)
51
51
(ksi)
66
67
(MPa)
454
461
(ksi)
59
61
(MPa)
406
420
(%)
0.90
0.96
ft.lbs.f
6
6
(J)
8
8
UTS
0.20% OFFSET
Elongation
Impact
2 w/o Cu
2 w/o Cu + V
Condition
Sintered
Sintered
2 w/o Cu
2 w/o Cu + V
Aged
Aged
7.00
6.98
52
52
66
71
454
489
63
67
433
461
0.62
0.82
5
5
7
7
2 w/o Cu
2 w/o Cu + V
HT2
HT
7.00
6.98
66
68
88
108
605
743
-----
-----
-----
-----
-----
1.
1
Aged at 538 oC (1000 oF) for 1h in nitrogen
2. Austenitiized at 900 oC (1650 oF) for 1h-oil quench and temper at 204 oC (400 oF)
Microstructures
The PM alloys were examined using light optical microscopy in an attempt to determine
if the differences in strength were related to the attendant microstructures. Figure 10
illustrates representative microstructures of the two alloys after aging. Both alloys are
composed primarily of lamellar pearlite with varying amounts of ferrite and a small
percentage of grain boundary carbides. The vanadium-containing alloy appears to
contain a slightly greater percentage of ferrite and more of the blocky, grain boundary
carbides than the vanadium-free alloy. The largest effect can be seen by comparing the
ferrite grain size of the two alloys. The vanadium-containing alloy is substantially finer
in grain size, indicating that there is inhibition to grain growth when compared with the
Fe- 2 w/o Cu.
(a)
(b)
Figure 10. Representative microstructures of Fe- 2 w/o copper in the aged condition: (a) vanadium-free
(b) with vanadium.
It is in the reaustenitized and oil quenched conditions that the most significant
microstructural differences can be seen. Figure 11 shows typical cross-sectional areas
from tensile bars fabricated from each alloy. Fe-2 w/o Cu, exhibits a microstructure of
plate martensite with a small amount of retained austenite. In contrast, the vanadiumcontaining alloy exhibits a lesser percentage of plate martensite and retained austenite,
but with a substantial amount of bainite. The finer grain size of the vanadium-containing
alloy is consistent with its higher strength; in addition, the bainite contributes to increased
ductility, in contrast to the plate martensite in the vanadium-free alloy.
(a)
(b)
Figure 11. Representative microstructures of Fe- 2 w/o copper after heat treatment: (a) vanadium-free (b)
with vanadium.
CONCLUSIONS
•
Copper in Fe-Cu PM alloys leads to precipitation and an increase in strength and
apparent hardness when aged at 538 oC (1000 oF).
•
Aging at 538 oC (1000 oF) of Fe-Cu PM alloys containing carbon leads to
tempering of the carbides and a reduction in strength and apparent hardness
•
Chromium, molybdenum and silicon stabilize the carbides at the precipitation
temperature for copper but the amount of each alloying element needed is high
and therefore costly.
•
The use of microalloying elements (< 0.20 w/o) such as niobium and vanadium
result in temper resistance in a Fe-Cu PM alloys at 538 oC (1000 oF) and provide a
cost effective way to increase the strength of alloys in the Fe-Cu system.
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Corporation, Newfield, New Jersey, 2000.